Blends of amorphous and semicrystalline polymers having shape memory properties

ABSTRACT

A shape memory polymeric material has a glass transition temperature, T g , exceeding room temperature and exhibits rubber modulus and elasticity derived substantially from physical crosslinks is prepared by blending components including one crystalline polymer and two amorphous polymers. Methods of preparing the polymeric materials and uses of the polymeric materials, for example, as smart medical devices, are also disclosed.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application is a division of U.S. patent application Ser. No.11/624,952, filed Jan. 19, 2007, which is a division of U.S. patentapplication Ser. No. 10/683,558, filed Oct. 10, 2003 and issued as U.S.Pat. No. 7,208,550 B2, which in turn claims the benefit of the followingprovisional patent applications: Ser. No. 60/418,023, filed Oct. 11,2002; Ser. No. 60/466,401, filed Apr. 29, 2003; and Ser. No. 60/488,323,filed Jul. 18, 2003. Each of the foregoing patent applications isincorporated herein by reference to the extent not inconsistent with thepresent disclosure.

TECHNICAL FIELD

The present disclosure relates to shape memory materials and, moreparticularly, to blends of an amorphous polymer, such as poly(vinylacetate) (PVAc), and a semicrystalline polymer, which can be for examplepoly(lactic acid) (PLA) or poly(vinylidene fluoride) (PVDF). The presentdisclosure also relates to methods for the preparation of such blendsdemonstrating shape memory effects and to applications thereof.

BACKGROUND OF THE DISCLOSURE

Shape memory materials feature an ability to transform shape from atemporary, frozen, shape to a permanent shape when triggered by anenvironmental stimulus, such as heat, light, or vapor. Used creatively,this phenomena can be exploited for a wide range of applications. Manypolymers intrinsically show shape memory effects, e.g., on the basis ofrubber elasticity, combined with vitrification or crystallization butwith varying characteristics such as strain recovery rate, workcapability during recovery, and retracted state stability. Among thefirst shape memory polymers (SMPs) reported as such was crosslinkedpolyethylene, which was discovered and patented in 1971 by RadiationAppliances, Inc., and a methacrylic acid ester reported by theVernon-Benshoff Co. for use as a denture material. The mechanism ofstrain recovery for such a material was identified as being fardifferent from that of shape memory alloys (SMAs), which are basedlargely on nickel-titanium alloys.

More particularly, a shape memory polymer is a super-elastic rubber,when the polymer is heated to a rubbery state, it can be deformed underresistance of ˜1 Mpa modulus, and when the temperature is decreasedbelow either a crystallization temperature or a glass transitiontemperature, the deformed shape is fixed by the lower temperaturerigidity while, at the same time, the mechanical energy expended on thematerial during deformation is stored. Thereafter, when the temperatureis raised above the transition temperature (T_(m) or T_(g)), the polymerwill recover to its original form as driven by the restoration ofnetwork chain conformational entropy. The properties of the SMPs will beclosely linked to the network architecture and to the sharpness of thetransition separating the rigid and rubber states. Compared with SMAs,SMPs have an advantage of (i) high strain, to several hundred percentbecause of the large rubbery compliance while the maximum strain of aSMA is less than 8%. (ii) facile tuning of transition temperaturesthrough variation of the polymer chemistry; and (iii) processing ease atlow cost.

A concurrently filed application for patent filed by inventors hereindiscloses the synthesis and characterization of thermally stimulatedSMPs having different thermomechanical properties and their use invarious applications, including as medical devices and mechanicalactuators. The disclosed materials span a range of room temperaturemoduli, from rigid glassy materials having storage moduli of several GPato compliant rubbers with moduli as low as tens of MPa. Moreover, theretracting (rubbery) moduli have been tuned over the range 0.5<E<10 MPa,as dictated by the end application. One example of such an SMP ischemically crosslinked polycyclooctene (PCO), a stiff semicrystallinerubber that is elastically deformed above T_(m) to a temporary shapethat is fixed by crystallization. Fast and complete recovery of grossdeformations are achieved by immersion in hot water. Other SMPs offeringtunable critical temperatures and rubber modulus have been synthesizedusing a thermosetting random copolymer formed from two vinyl monomersthat yield controlled T_(g) and casting-type processing. Such copolymerswere crosslinked with a difunctional vinyl monomer as crosslinker, theconcentration of crosslinker controlling the rubber modulus and thus thework potential during recovery. In addition to their shape memoryeffects, these materials are also castable allowing for processing morecomplex shapes and they are optically transparent. The use of chemicalcrosslinking in both of these cases, however, limits the types ofprocessing possible and permanently sets the equilibrium shape at thepoint of network formation.

Semicrystalline thermoplastic polymers with sharp T_(g)> roomtemperature and low crystallinity are also good candidates for shapememory, while offering the advantage of melt processing above T_(m),allowing repeated resetting of the equilibrium shape by relaxing stressin the fluid state. Representative of the polymers in this class of SMPsare polyurethanes whose soft domains are semicrystalline with lowmelting points (but higher than T_(m)) and whose hard domains feature ahigher melting point only exceeded during processing. The effect of thehard segments and soft segments on the shape memory effects have beeninvestigated. In addition to such polyurethanes, block copolymers ofpolyethylene terephthalate (PET) have been synthesized for their shapememory effects.

Miscible blends of a semicrystalline polymer with an amorphous polymerhave been investigated in the past, though not with respect to shapememory behavior, due to their attractive crystalline properties andmechanical properties. For those blends that are miscible at themolecular level, a single glass transition results, without broadening,an aspect important to shape memory. Additionally, in such miscibleblends the equilibrium crystallinity (which controls the plateau modulusbetween T_(g) and T_(m) where shape fixing is performed) also changesdramatically and systematically with blend composition; i.e., relativelevels of each component. Although numerous blends of this type havebeen investigated, there has been no known disclosure of the utilizationof such blends for their shape memory properties.

SUMMARY OF THE DISCLOSURE

In accordance with the present disclosure, SMPs with relatively highmodulus in the fixed state at room temperature, having a tunable andsharp transition, the permanent shape of which can be remoldedrepeatedly above certain melting temperatures, are prepared by theblending or mixing of a crystalline polymer (C) with an amorphouspolymer (A), such that they are a single miscible phase in the moltenstate (allowing processing to stress-free native states), butcrystalline to a limited and tailored extent and which further vitrifyon cooling to room temperature. Examples for (C) include, but are notlimited to, poly(vinylidene fluoride) (PVDF) (T_(g)=−35° C., T_(m)=175°C.), polylactide (PLA) (T_(g)=56° C., T_(m)=165° C.), and its copolymerssuch as poly(L-lactide), poly(D,L-lactide), poly(lactide-co-glycolide),poly(lactide-co-caprolactone), poly(lactide-co-1,5-dioxepan-2-one),poly(lactide-co-trimethylene carbonate), polyglycolide, poly(3-hydroxybutyrate) and its copolymers, polyanhydrides, poly(ethylene glycol)(PEG), polyethylene, polyethylene-co-vinyl acetate, poly(vinyl chloride)(PVC), and poly(vinylidene chloride) (PVDC) and copolymers of vinylidenechloride and vinyl chloride. Examples for (A) include poly(vinylacetate) (PVAc) (T_(g)=35° C.), poly(methyl acrylate) (PMA), poly(ethylacrylate) (PEA), atactic poly(methyl methacrylate) (aPMMA), isotacticpoly(methyl methacrylate) (iPMMA), syndiotactic poly(methylmethacrylate) (sPMMA), and other poly(alkyl methacrylates).

Plasticizers may be included in the disclosed blends according to thepresent disclosure. Exemplary plasticizers include, but are not limitedto, di-n-butyl phthalate (DBP), di-n-octyl phthalate (DNOP),di(2-ethylhexyl)phthalate (DOP), di-2-ethylhexyl) isophthalate (DOIP),bis-(2-ethylhexyl)terephthalate (DOTP), di-n-decyl phthalate (DNDP),di-cyclohexyl phthalate (DCHP), di-octylsebacate (DOS), Hexamoll® DINCH,EVA, EVA-carbon monoxide terpolymer (Elvaloy), poly(alkylenealkanoates). DOP can also be used to decrease the glass transition ofPVC.

The disclosed polymer blends may also be compounded to include a finelydivided particulate material, such as clay, silica or TiO₂.

For high stiffness in the temporary form, a need exists for shape memorypolymers having a T_(g) greater than room temperature, but with atailored rubber modulus and elasticity derived from physical rather thanchemical crosslinks.

Preferably, exemplary SMPs of the present disclosure are achieved byblending or mixing amorphous poly(vinyl acetate) (PVAc) (T_(g)=35° C.)with semicrystalline polylactide (PLA) (T_(g)=56° C., T_(m)=165° C.) orpoly(vinylidene fluoride) (PVDF). The disclosed polymers show completemiscibility at all blending ratios with a single glass transitiontemperature, while crystallization (exclusive of PVAc) is partiallymaintained. The T_(g)'s of the blends are employed as the criticaltemperature for triggering the shape recovery while the crystallinephases function as physical crosslinking sites for elastic deformationabove T_(g), but below T_(m). Shape memory polymers have receivedincreased attention recently, prompted by an expanding range ofpotential end-use applications, especially for development of biomedicalengineering tools and as medical devices.

Presently preferred polymer blends according to the present disclosureare formed from poly vinyl acetate(PVAC) and poly(lactic acid) (PLA) orpoly(vinylidene fluoride) (PVDF). However, examples of other suitableblends include the pair PVDF/PMMA and ternary blends of PVDF/PMMA/PVAc.The PMMA and the combination of PMMA/PVAc serve the same role as PVAc inthe blends as has been previously described. An advantage of adding PMMAis that the critical temperature can be increased arbitrarily to about80° C. and the room temperature modulus can also be increased. The PVDFmay be substituted by poly(vinylidene chloride) (PVDC), by copolymers ofpoly(vinylidene chloride/poly(vinyl chloride), or by any “C” polymervide supra.

It has further been found that blending or mixing poly(vinyl chloride)with poly(butyl acrylate) or poly(butyl methacrylate) (PVC/PBA) hascertain advantages. In the PVDF/PVAc case, PVAc simultaneously lowersthe crystallinity of PVDF while increasing the T_(g). PVC may serve thesame role as PVDF, but it already has a low degree of crystallinity, buta relatively high T_(g) (80° C.). Thus, in this exemplary embodiment ofthe present disclosure, the second component (PBA) serves only thefunction or role of decreasing T_(g). This function/role can also beachieved with small molecule plasticizers, most notably dioctylphthalate(DOP), but it is presently preferred to use a biocompatible polymericplasticizer for intended implantable medical applications. The range ofPBA compositions is 10-40%, with 20% being the most advantageous,yielding a T_(g)˜40° C.

Further advantageous features and functions associated with the polymerblends of the present disclosure will be apparent from the detaileddescription which follows.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 illustrates a comparison of the DSC traces of PLA (top) quenchedfrom T=180° C. or annealed at T=110° C. for one (1) hour (bottom);

FIG. 2 illustrates DSC traces for PLA/PVAc blends following annealingfor one (1) hour at T=110° C.;

FIG. 3 illustrates glass transition temperatures as measured followingquenching of the PLA/PVAc blends;

FIG. 4 illustrates tensile storage modulus versus temperature for arange of PLA/PVAc; and

FIG. 5 illustrates shape recovery of coiled PLA/PVAc (30:70) blend onexposure to water at T=65° C.

DETAILED DESCRIPTION OF EXEMPLARY EMBODIMENTS

Miscible blends of the present disclosure show or exhibit sharp andtunable transition temperatures and adjustable stiffness above theirtransition temperatures. The disclosed blends also show/exhibitexcellent shape recovery effect with the recovery temperature andretracting force that is adjustable according to the transitiontemperature of the amorphous polymer and the degree of crystallinity ofthe homopolymers, respectively. The shape recovery can be finishedwithin seconds when heated to an appropriate level above the transitiontemperature, e.g., 20° C. above the transition temperature. Additionaladvantages associated with the materials of the present disclosureinclude that the materials are rigid at room temperature, the polymersgenerally are biocompatible and can be used as medical devices andimplants, and the products also can be dyed to any color or renderedradio-opaque for x-ray radiography according to applicationrequirements.

The shape recovery temperatures of particular polymer blends accordingto the present disclosure depend on the glass transition of the blends.Theoretically, the blended polymer will recover above the glasstransition temperature and below the melting temperature. Temperaturesthat are on the order of 20° C. above the glass transition temperatureare preferably used to make the shape recovery fast. More preferably atemperature in the rubbery plateau region for the particular polymerblend is selected in order to have a fast recovery and a predictable anddesired retracting force.

Blends of the present disclosure may be successfully used in connectionwith a wide variety of applications including, without limitation, thefollowing applications:

a) Stents, patches and other implants for human health care.

b) Surgical tools requiring adjustable shape but high stiffness.

c) Arbitrarily shape-adjustable structural implements, includingpersonal care items (dinnerware, brushes, etc.) and hardware toolhandles.

d) Self healing plastics.

e) Medical devices.

f) Impression material for molding, duplication, rapid protyping,dentistry, and figure-printing.

g) Toys.

h) Reversible embossing for information storage.

i) Temperature sensors

j) Safety valves

k) Heat shrink tapes or seals

l) Heat controlled couplings and fasteners.

m) Large strain, large force actuators

EXAMPLES

To further illustrate advantageous features, functions and uses ofexemplary polymer blends according to the present disclosure, thefollowing non-limiting examples are provided. As will be readilyapparent to persons of skill in the art, the following examples aremerely illustrative of aspects of the present disclosure, and should notbe viewed as absolute and thus should not be considered to be limitingwith respect to potential polymeric materials, processing conditions(e.g., relative percentages, temperatures and time periods) and/orend-use applications that may be advantageously employed according tothe present disclosure. Physical properties and processing conditionsset forth in the following examples are merely illustrative of suchproperties/conditions, and should not be viewed as limiting the scope orutility of the present disclosure.

Materials

Poly(vinyl acetate) having M_(w)=500,000 g/mol was purchased fromAldrich and vacuum-treated at room temperature for 2 days to eliminatevolatiles, including moisture. The poly(lactide) having M_(w)=130,000g/mol was provided by Prof. S. Huang of the University of Connecticutand was dried at T=60° C. overnight before use. The material was foundto comprise 92% L-isomer and 8% D-isomer. Poly(vinylidene fluoride)having a M_(W)=180,000 g/mol was also purchased from Aldrich andvacuum-treated at room temperature for 2 days to eliminate volatiles,including moisture.

Processing and Fixing the Primary Shapes

Melt blending of PLA/PVAc and PVDF/PVAc using varying blend ratios wasperformed in a 30 ml Brabender twin-screw mixer. The mixer wasequilibrated at T=180° C. for 5 minutes after which the mixer bladerotation was adjusted to 25 RPM and the premixed polymer pellets addedto the chamber over the course of 1 minute. The polymers were mixed for10 minutes to ensure good dispersion. Nitrogen was purged through thechamber to mitigate potential oxidative degradation during mixing. Aftermixing, the blend was taken out of the chamber, cooled to roomtemperature, and then pressed between heated platens of a Carver pressat 180° C. for 5 minutes under a load of 8 metric tons. A spacer wasused to control the thickness of the film and rapid cooling to roomtemperature was carried out. The films thus formed were used for thermaland mechanical characterization.

Characterizations

The thermal properties (thermogravimetric analysis (TGA)) of thepolymers were measured preceding differential scanning calorimetry (DSC)measurements and melt blending to ensure that no solvent vapor waspresent in the polymers. For this purpose, a TA Instruments (TA 2950)TGA was used with samples of the polymers heated from 25° C. to 600° C.at a heating rate of 20° C./minute in a nitrogen atmosphere. APerkin-Elmer DSC-7 was used for DSC analysis. Ten (10) mg samplesprepared from pellets and heat pressed films were first heated from 20°C. to 200° C. at a heating rate of 10° C./minute (1^(st) heat), thencooled to 20° C. at the same rate, and finally the samples were reheatedto 200° C. at a heating rate of 10° C./minute (2^(nd) heat). The T_(g)of each blend was determined from the midpoint of heat capacity changeobserved during 2^(nd) heating.

Wide-angle x-ray scattering (WAXS) analysis was performed using a BRUKERGADDS-4 instrument with a Cr source of radiation (λ=2.291 Å) and intransmission mode. The voltage and the current used were 40 kV and 40mA, respectively, and the exposure time was 30 minutes. The scatteringpatterns were collected on a HiStar area detector with asample-to-detector distance of 6.0 cm. Intensity profiles (I versus 2θ)were determined by azimuthal averaging at each 2θ position of theisotropic patterns. The data were then analyzed with Peakfit™ software(SPSS Science) to find the peak positions and the relative intensity ofeach peak.

The storage and loss tensile moduli of the blends were measured bydynamic mechanical thermal analysis (DMTA) in tensile mode using TAInstruments DMA2980. The method used followed a temperature rampspanning −100° C.<T<150° C. at a heating rate of 4° C./minute, a fixedoscillation frequency of 1 Hz, and a maximum tensile strain of about0.5%. The sample geometries were rectangular bars of dimensions 10×2×1mm.

Isothermal stress-free shape recovery tests were carried out usingsamples cut into a rectangular shape, stained for optical contrast, andthen heated to T=65° C. for bending into a temporary helical shape. Thedeformed sample was fixed by cooling in ice water, thus vitrifying thesample. The resulting deformed sample was subsequently dropped into awarm water bath at a predetermined temperature and the shape recoverymonitored visually with a video camera and digital frame-grabbercollecting images at a rate of 20 frames-per-second.

Results

The TGA results demonstrated that both PLA and PVAc are stable forT<300° C. Above this temperature, PLA degrades completely (no charyield), while the PVAc degrades to yield an intermediate char yield of25 wt % for 375<T<425° C., but complete degradation occurs above 450° C.Blend processing and thermal and dynamic mechanical analyses (DSC andDMA) were performed below 250° C., to completely avoid degradation.

The crystallization behavior of semicrystalline PLA was investigated viaDSC. The PLA samples were first heat pressed at 180° C. for 10 minutesand then quenched to room temperature with water cooling. One sample wasdirectly analyzed by DSC, while another was first annealed at 110° C.(=½(T_(g)+T_(m))) for 1 hour to reach an equilibrium level ofcrystallinity. FIG. 1 shows a comparison of thermal behavior for thesetwo samples. It was observed that quenching the PLA melt results in alow degree of crystallinity and virtually no recrystallization onheating, both indicating slow crystallization. Annealing at 110° C. for1 hour results in significant crystallization evidenced by a largemelting endotherm at T=155° C. The melting temperature did not shiftdramatically due to annealing, but the endotherm shape did change.Complementary WAXD experiments yielded the same conclusions.

The crystallization behavior of polymer blends was also analyzed. All ofthe samples were heat pressed at 180° C. for 10 minutes and thenannealed at 110° C. for 1 hour before thermal analysis, providing astandard condition for extensive crystallization. FIG. 2 summarizes thefirst DSC heating trace of the samples measured after annealing. Theresults indicate that PVAc itself is amorphous (though with largephysical aging content), but that incorporation of PLA leads tocrystallization in proportion to the PLA wt-%. Also, the peak endothermpositions (melting transitions) shift slightly to higher temperatureswith increasing PLA content. Quenching these samples to T=20° C. andreheating to 200° C. again showed clearly that single T_(g)'s areobserved and that crystallization can be largely suppressed. Importantlyfor shape memory, the single glass transition events were not broadenedin the blends relative to the pure components, suggesting that theamorphous phase was quite homogeneous for all of the blends. Theobserved T_(g) values are plotted in FIG. 3 along with the best fit withthe Fox equation, showing slight positive deviation. This leads to aconclusion that strong interactions between the two polymers thatreduces free volume of the polymer blends and hence, increased glasstransition temperature relative to the Fox equation prediction, hasoccurred.

In order to elucidate the effect of PVAc on the degree of crystallinityand the crystal structures, the crystalline diffraction patterns wereobserved via wide-angle x-ray diffraction. The results indicate that thePVAc phase has only an amorphous halo, thus being totally amorphous,while the PLA exhibits three very strong diffraction peaks at 2θ=22.3°,25.0° and 28.6°, corresponding to d-spacings of 5.92, 5.29, and 4.64 A°,respectively. Upon addition of PVAc, all of the peak intensities weredepressed, but the peak positions remained essentially unchanged.Consistent with the DSC results, the degree of crystallinity increasesin proportion to PLA addition. From the peak width at half height, itwas found that the crystalline lamellae size did not decrease, althoughthe degree of crystallinity decreased, with increasing PVAc content.This means that the decrease in crystallinity and depression of themelting transitions are not due to a change of crystal size, but rathermay be due to a thinning of the lamellae thickness or to a decrease ofthe crystal concentrations.

The storage modulus of the polymer blends was also measured using DMTA,first investigating the effects of annealing on the storage modulus.Below their glass transition temperatures, T_(g), both samples exhibitsimilar high storage moduli (3 GPa), as well as similar softeningpoints. When heated above T_(g), the storage modulus of thermallyquenched samples decreases sharply to about 2 MPa. However, furtherincreasing the temperature induces a modulus increase attributed torecrystallization of the samples at higher temperatures. This alsoproved that the sample is not in an equilibrium state and that itsmechanical properties in the rubbery region depend on thermal history.To reach equilibrium, the sample was annealed at 110° C. for 1 hour aspreviously described for DSC analyses. The storage modulus above T_(g)shifts to about 200 MPa until melting, the increase being due to anincrease of the degree of crystallinity on annealing. In order to tunethe rubbery modulus at equilibrium state, PLA was blended in differentproportions to PVAc and annealed as above. Storage moduli for suchblends were measured and the results are plotted in FIG. 4. It can beseen that, below T_(g), all of the samples show similar large moduliwhile above T_(g) the moduli decrease to a plateau whose magnitudedepends on crystallinity and thus PLA content. This trend is inaccordance with that of DSC and XRD, and may be explained by the factthat the increase of storage moduli came from the physical crosslinkingformed by crystals and the filler effect of the high modulus crystallinephase.

Stress-free shape memory tests were carried out in hot water at 65° C.,with an annealed sample composed of 30% PLA. The real time shaperecovery was videotaped and selected images are shown in FIG. 5. Theresults show that the sample features quick and complete shape memorybehavior: the sample recovers to the original shape (straight bar)within 10 seconds, with most of the recovery being accomplished withinthe first several seconds.

The same characterizations were carried out on the blends of PVDF andPVAc as above described. The TGA and DSC results show that PVDF is alsothermally stable up to 300° C., and the mixtures form only one glasstransition, the values fall between the Tgs of the two homopolymers andchanges with changing composition. At the same time, the melting pointsand the degrees of crystallinity were depressed with the incorporationof amorphous PVAc.

The storage moduli of the blends, which give the rigidity of thematerials, were also measured. The results are similar to those of thePLA/PVAc blends, the PVDF/PVAc blends being very rigid below thecritical temperatures (T_(g)), and featuring sharp modulus changes atthe Tg to a plateau modulus ranging from several MPa to tens of MPa,depending on the degree of crystallinity of the blends. These plateaumoduli can be tuned by adjusting the degree of crystallinity of theblend, that is, by adjusting the blend composition.

In summary, shape memory polymers were obtained by blending asemicrystalline polymer, for example, PLA, with an amorphous polymer,for example PVAc. The polymer blends are totally miscible at all blendratios within the experimental ranges and form only one single glasstransition temperature for each formulation. Additionally, the degree ofcrystallinity of the blends decreases monotonically with increasing PVAcand PVAc and PVDF fraction. This, in turn, governs the rubbery modulusimportant to shape memory.

Thus, the present disclosure advantageously provides a shape memorypolymeric material that is characterized by a T_(g) exceeding roomtemperature whose rubber modulus and elasticity are derivedsubstantially from physical crosslinks comprising a blend of acrystalline polymer selected from the group consisting ofpoly(vinylidene fluoride), polyglycolides, polylactide and copolymersthereof, poly(hydroxybutyrate), poly(ethylene glycol), polyethylene,polyethylene-co-vinyl acetate, poly(vinyl chloride), poly(vinylidenechloride) and copolymers of poly vinylidene chloride and poly vinylchloride with an amorphous polymer selected from the group consisting ofpoly(vinyl acetate), poly(methyl acrylate), poly(ethyl acrylate),atactic poly(methyl methacrylate), isotactic poly(methyl methacrylate)and syndiotactic poly(methyl methacrylate).

The present disclosure also advantageously provides a method ofpreparing a shape memory polymer material characterized by a T_(g)exceeding room temperature whose rubber modulus and elasticity isderived substantially from physical crosslinking comprising meltblending a crystalline polymer selected from the group consisting ofpoly(vinylidene fluoride), polylactide, poly(hydroxybutyrate),poly(ethylene glycol), polyethylene, polyethylene-co-vinyl acetate,poly(vinyl chloride), poly(vinylidene chloride) and copolymers of polyvinylidene chloride and poly vinyl chloride with an amorphous polymerselected from the group consisting of poly(vinyl acetate), poly methylacrylate, poly ethyl acrylate, atactic poly methyl methacrylate,isotactic poly methyl methacrylate and syndiotactic poly methylmethacrylate at a temperature of 10˜20° C. above the melting temperatureof the crystalline polymers, for a time sufficient to ensure goodmixing, cooling the resultant blend to room temperature, introducingsaid blend into a press maintained at about 180° C., applying pressureto said blend and then rapidly cooling the film thereby formed to anannealing temperature T_(g)<T_(a)<T_(m), where it is held untilcrystallization is complete and following which the film is cooled toroom temperature.

Although the polymer blends and processing methodologies of the presentdisclosure have been described with reference to specific exemplaryembodiments thereof, the present disclosure is not to be limited to suchexemplary embodiments. Rather, as will be readily apparent to personsskilled in the art, the teachings of the present disclosure aresusceptible to many implementations and/or applications, withoutdeparting from either the spirit or the scope of the present disclosure.Indeed, modifications and/or changes in the selection of specificpolymers, polymer ratios, processing conditions, and end-useapplications are contemplated hereby, and such modifications and/orchanges are encompassed within the scope of the present invention as setforth by the claims which follow.

1. A shape memory polymeric material characterized by a T_(g) exceedingroom temperature whose rubber modulus and elasticity are derivedsubstantially from physical crosslinks comprising a ternary blend of onecrystalline polymer and two amorphous polymers; wherein the shape memorypolymeric material has a melting temperature, T_(m), and a glasstransition temperature, T_(g), exceeding room temperature; and whereinthe shape memory polymeric material has rubber modulus and elasticityderived substantially from physical crosslinks formed by annealing theblend at a temperature greater than T_(g) and less than T_(m).
 2. Theshape memory polymeric material of claim 1, wherein the one crystallinepolymer is selected from the group consisting of poly(vinylidenefluoride), polyglycolides, polylactides, poly(lactide-co-glycolide),poly(lactide-co-caprolactone), poly(lactide-co-1,5-dioxepan-2-one),poly(lactide-co-trimethylene carbonate), poly(hydroxybutyrate),poly(ethylene glycol), polyethylene, poly(ethylene-co-vinyl acetate),poly(vinyl chloride), poly(vinylidene chloride) and copolymers ofvinylidene chloride and vinyl chloride.
 3. The shape memory polymericmaterial of claim 1, wherein the two amorphous polymers are eachselected from the group consisting of poly(vinyl acetate), poly(methylacrylate), poly(ethyl acrylate), and atactic poly(methyl methacrylate).4. The shape memory polymeric material of claim 1, wherein the ternaryblend comprises amorphous poly(methyl methacrylate); amorphouspoly(vinyl acetate); and a crystalline polymer selected from the groupconsisting of poly(vinylidene fluoride), poly(vinylidene chloride), andpoly(vinylidene chloride-co-vinyl chloride).
 5. The shape memorypolymeric material of claim 1, wherein the ternary blend comprisescrystalline poly(vinylidene fluoride), amorphous poly(methylmethacrylate), and amorphous poly(vinyl acetate).
 6. The shape memorypolymeric material of claim 1, wherein the ternary blend comprisescrystalline poly(vinylidene chloride), amorphous poly(methylmethacrylate), and amorphous poly(vinyl acetate).
 7. The shape memorypolymeric material of claim 1, wherein the ternary blend comprisescrystalline poly(vinylidene chloride-co-vinyl chloride), amorphouspoly(methyl methacrylate), and amorphous poly(vinyl acetate).
 8. Theshape memory polymeric material of claim 1, wherein the physicalcrosslinks are formed by annealing the blend at an annealing temperaturegreater than the glass transition temperature of the blend and less thanthe melting temperature of the blend, wherein the annealing is conductedfor a time effective to produce equilibrium crystallinity.
 9. The shapememory polymeric material of claim 1, wherein the one crystallinepolymer is a polyanhydride.
 10. The shape memory polymeric material ofclaim 1, wherein the one crystalline polymer is a polyester.